Turk J Phys
() : 1 – 12
c⃝ TUBITAK
doi:10.3906/fiz-1412-7
Turkish Journal of Physics
http :// journa l s . tub i tak .gov . t r/phys i c s/
Research Article
Synthesis and properties of a spinel cathode material for lithium ion battery with
flat potential plateau
Ahmed Abdulrahman Ahmed AL-TABBAKH1,∗, Norlida KAMARULZAMAN2,Aseel Basim AL-ZUBAIDI3
1Department of Physics, Al-Nahrain University, Jadiriya, Baghdad, Iraq2Centre for Nanomaterials Research, Institute of Science, Universiti Teknologi MARA,
Shah Alam, Selangor, Malaysia3Department of Materials Engineering, University of Technology, Baghdad, Iraq
Received: 18.12.2014 • Accepted/Published Online: 03.04.2015 • Printed: ..201
Abstract: A potential cathode material for lithium ion battery was synthesised by combustion reaction. The thermal
behaviour of the as-synthesised precursor was measured using a thermogravimetric analyser and the range of calcination
temperature from 500 ◦C to 800 ◦C was determined. X-ray diffraction analysis showed that all calcined powders
crystallised in the cubic spinel structure of the Fd3m space group. The particle size distributions and morphologies
of the powders were obtained using a particle size analyser and scanning electron microscope. The effect of calcination
temperature on the electrochemical performance was investigated using galvanostatic charge/discharge measurements.
The potential profiles of charge and discharge exhibited flat plateaus, emphasising that the synthesised batteries are
not super capacitors. Results showed that the batteries prepared from powders calcined at 800 ◦C exhibited higher
coulombic efficiency and better cycling performance than powders calcined at 500 ◦C.
Key words: Li-ion batteries, transition-metal compounds, combustion synthesis, electrochemical properties, X-ray
diffraction
1. Introduction
The profile of the charge/discharge potential curve with capacitor-like behaviour is one major weakness observed
in lithium ion batteries. The capacitor-like behaviour is the linear variation of the battery potential in accordance
with the intercalation/deintercalation of lithium ions into/from the active material. It is attributed to the low
electronic conductivity and low lithium ion diffusion coefficient of the active material and is most likely to
increase with the plateau voltage and the charge/discharge currents [1]. The usefulness of the battery material
for industrial and technological purposes depends largely on moderating this problem. Lithiated transition
metal oxides of spinel structures have been recognised as potential cathode materials for high energy and high
power density lithium secondary batteries due to high conductivity and high lithium diffusion coefficient [2–
4]. Cathode materials containing manganese and nickel in the form of LiNixMn2−xO4 have been extensively
investigated due to their high operating potentials, lower toxicity than cobalt-containing cathodes, and low
cost [5,6]. However, increasing the nickel content of the cathode material invites disadvantages like formation
of impurities during synthesis and fast capacity fade at elevated temperatures [7,8]. Partial substitution of
either or both of the aforementioned transition metals with other cations like Fe is thought to overcome these
∗Correspondence: [email protected]
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AL-TABBAKH et al./Turk J Phys
disadvantages and improve the electrochemical performances of the cathodes [9]. However, the high voltage
potential of these materials is always accompanied by noticeable potential drop across the high potential
plateaus. In this work we report the synthesis, characterisation, and properties of a spinel cathode material
of the form Li1.03Mn1.72Ni0.16Fe0.04O4 having a flat potential plateau with minimal potential drop. The
effect of calcination temperature on the electrochemical performance of the material was investigated using
galvanostatic charge/discharge measurements. The results reported in the present discussion suggest that the
present material is a potential candidate for lithium rechargeable batteries for high power density applications.
2. Experimental work
2.1. Synthesis and characterisation of the material
Synthesis of the cathode material was achieved by a self-propagating combustion reaction [10,11]. Stoichiometric
amounts of LiNO3 , Mn(NO3)2 .4H2O, Ni(NO3)2 .6H2O, and Fe(NO3)3 .9H2O were dissolved in deionised water
at room temperature. The solution was then kept on magnetic stirring for 3 h to confirm dissolution of the
nitrates. An aqueous solution of citric acid (C6H8O7), used as fuel for the combustion reaction, was added to
the homogenised mixture at 100 ◦C. Fuel amount was determined from the oxidation to reduction ratio of the
reactants. Prior to the combustion, excess water was evaporated from the solution at a moderate temperature
(below 150 ◦C) until a dense solution was obtained. The temperature of the solution was finally raised at a
high rate to the combustion temperature to prevent formation of impurities [12]. The combustion reaction was
a smouldering (flameless) type and self-propagating, resulting in a fine powdered precursor. Figure 1 shows the
phases of the synthesis reaction just before the combustion (i.e. Figure 1a), during combustion (Figure 1b), and
just after the combustion took place (Figure 1c). In Figure 1b it is clear that combustion took place consistently
throughout the reaction vessel.
(a) (b) (c)
Figure 1. The phases of the combustion reaction: (a) formation of frothed solution of the reactants, (b) instant of
combustion, and (c) formation of precursor just after combustion of the reactants.
The produced precursor from the reaction was thermally analysed using a high-performance modular
thermogravimetric SETSYS Evolution Analyser from SETARAM Instrumentation to estimate the best calcina-
tion temperature. A few milligrams of the precursor were loaded on the analyser crucible. Thermogravimetric
(TG) and heat flow graphs were obtained at a 10 ◦C/min heating rate under the flow of a protective gas.
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AL-TABBAKH et al./Turk J Phys
Calcination of the precursors was achieved by admitting the precursor powder into a hot furnace at a pre-set
temperature to prevent formation of oxide impurities and to produce pure powders [12]. The calcined powders
were then processed by a high-energy ball mill operated at 600 rpm using 9 balls of 4 g each to produce finer
powders. It is well known that size reduction of the material particles may enhance the material’s attainable
capacity and rate capability due to increase of surface to volume ratio [13].
The powders were then characterised by particle size distribution and X-ray diffraction (XRD) measure-
ments. The size distributions were measured using a Malvern Mastersizer Range particle size analyser from
Malvern Instruments [14]. Precise amounts of the powders were dispersed in distilled water and sonicated
for 15 min to achieve the required dispersion of the particles. Measurements were repeated several times to
ensure consistency of results. The XRD measurements were carried out with a PANalytical X’pert Pro MDP
instrument with a real-time multiple strip type solid-state detector (known as the X’celerator), spinner holder,
and Bragg–Brentano reflection-parafocusing geometry. The morphologies of the powders were examined using a
JSM-6700F field emission scanning electron microscope (SEM) from JEOL. A few randomly selected specimens
were investigated under similar working conditions and high magnification.
2.2. Cathode fabrication, battery assembly, and test
Slurries were formed and homogenised from a mixture of 0.025 g of the active material, 0.003 g of activated
carbon, 0.0027 g of polyvinylidene fluoride (PVDF) binder, and a few droplets of acetone. These slurries were
then pasted on metallic mesh serving as current collectors. The fabricated cathodes were dried at 200 ◦C for 24 h
before being admitted to the batteries’ fabrication chamber. The lithium ion batteries were assembled in Teflon
battery holders inside a glove-box from Unilab Analyse under a dry argon atmosphere. The assembled cells
were composed of the fabricated cathodes as the positive electrode, a porous polypropylene separator (Celgard
2500), and a Li foil anode, all immersed in 1 M lithium hexaflurophosphate LiPF6 electrolyte in a 1:1 volume
ratio of ethylene carbonate and dimethyl carbonate. Each electrode had dimensions of (1×1) cm2 . The charge
and discharge equilibrium potential profiles were measured in the voltage ranges from open-circuit voltage to
4.2 V and from 4.2 V to 1.0 V. The charge and discharge currents were fixed at 0.5 mA (20 mA/g). The
assembled batteries were operated for 20 cycles of charge and discharge at a constant current of 1 mA using a
16-channel automatic battery cycler WBCS3000 fromWonATech. Specific capacity values were determined from
the applied current, time of discharge, and active material loading in the cathodes (i.e. 0.025 g). Measurements
were achieved at room temperature and repeated several times to ensure reproducibility of results.
3. Results and discussion
The reaction is proposed to take place in three ideal steps, as shown schematically in Figure 2. In the first step
the metal nitrates react with water-producing metals, hydroxides, and nitric acid. The second step, where the
citric acid added, results in the formation of intermediate compounds of the metal carboxyl. The carboxyls are
expected to react with the HNO3 (formed in the first step) during the combustion of the mixture and produce
the final precursor, which is composed of the lithiated metal oxide (i.e. Li1.03Mn1.72Ni0.16Fe0.04O4), carbon
monoxide, nitrogen, and water. The three steps mentioned above are meant to take place in the ideal and
complete reaction, which is practically difficult to achieve. Therefore, some impurities of other compounds may
form during the synthesis procedure. FeMnO3 is one such impurity that has been identified from the XRD
patterns of the as-synthesised powders. This will be further discussed in the following paragraphs. Formation
of this impurity is expected to take place during the last step, where FeHCOOH reacts with MnHCOOH in
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AL-TABBAKH et al./Turk J Phys
the presence of nitric acid according to the formula (i.e. FeHCOOH + MnHCOOH + 2HNO3 → FeMnO3 +
3H2O + 2CO2 + N2). Formation of Fe2O3 is also possible as the temperature is raised above 200 ◦C and just
before the combustion of the mixture occurs. Nevertheless, Fe2O3 may react with the citric acid producing
FeHCOOH in the second step of the reaction in what is thought to be a self-correcting step (referred to with
the vertical dashed arrows in Figure 2). It is therefore thought that raising the temperature of the nitrates’
aqueous solution in the first step of the reaction (before adding C6H8O7 to the mixture) should not exceed
the formation temperature of the Fe2O3 . Other hydrocarbon impurities (e.g., LiC3 , LiC4 , LiC6 , and Li2C2)
may also form in the second step of the reaction, as shown in Figure 2 (horizontal dashed arrow). These are
expected to react with HNO3 to produce lithium hydroxide, which ultimately takes part in the third step of
the reaction (i.e. the combustion of the materials) [10,11,15].
Fe 2O3 LiNO3 , Mn(NO 3)2,
Ni(NO3)2, Fe(NO 3)3
H2O
HNO3 LiOH, Mn(OH) 2, Ni(OH)2,
Fe(OH) 3
C6H8O7 H2O + CO 2
Li 2C2, LiC 3, LiC 4, …
LiOH, MnHCOOH, NiHCOOH,
FeHCOOH
Li 1.03Mn1.72Ni0.16Fe 0.04O4 + FeMnO 3 + H2O + N2 + CO2
200 °C
Figure 2. The proposed reaction route for the synthesis of the cathode material by combustion method. The dashed
arrows indicate incomplete combustion and formation of impurity phases.
The TG and associate heat flow curves are shown in Figure 3. The TG curve exhibits four distinct regions
of mass losses as indicated by m1 through m5 in Figure 3. Based on the temperatures ranges and the nature
of the heat flow curve, the first region (m1 –m2) was attributed to loss of water. The second region (m2–m3)
is due to decomposition of some residual hydrocarbons that might not have taken part in the reaction. The
third region (from m3 to m4) was attributed to decomposition of FeMnO3 and formation of the pure phase of
spinel Li1.03Mn1.72Ni0.16Fe0.04O4 compound. The mass loss in the last region (m5) could only be attributed
to delithiation of the spinel Li1.03Mn1.72Ni0.16Fe0.04O4 by lithium evaporation. Mass losses, corresponding
masses, and molar masses of each region are presented in Table 1.
This analysis suggests that the as-synthesised precursor was composed of 94.5% Li1.03Mn1.72Ni0.16Fe0.04
O4 , 2.2% FeMnO3 , 1.7% hydrocarbons, and 1.6% residual water vapour. Based on the TG analysis the as-
synthesised precursor was divided into two quantities, one of which was calcined at 500 ◦C for 24 h and the other
at 800 ◦C for 24 h. Both powders were then subjected to ball milling to produce fine powders labelled hereafter
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AL-TABBAKH et al./Turk J Phys
TG
%
m 5
Heat Flow μV
m 4
Hea
t "
ow
( μ
V)
m1
m 2
m 3
Temperature (°C)
2
4
6
8
10
12
14
16
0 0
200
400 600
800
0
–2
–4
–6
–8
–10
Figure 3. TG and associate heat flow curves of the thermally analysed precursor. Masses and corresponding molecular
masses at m1 : 14 mg, 187.582 g/mol. m2 : 13.776 mg, 184.581 g/mol. m3 : 13.538 mg, 181.392 g/mol. m4 : 13.23 mg,
177.265 g/mol. m5 : 13.104 mg, 175.577 g/mol.
Table 1. Mass losses, molar masses, and composition of material at TG curve regions.
Region T (◦C) Mass (mg) Molar mass (g/mol) Mass loss (mg) Composition
m1 25 14.00 187.582 0.0
1Li1.03Mn1.72Ni0.16Fe0.04O4 +0.16747H2O +0.0503(hydrocarbons) +0.026FeMnO3
m2 166 13.775 184.581 0.224
Li1.03Mn1.72Ni0.16Fe0.04O4 +0.026FeMnO3 +0.0503(hydrocarbons)
m3 340 13.538 181.392 0.238Li1.03Mn1.72Ni0.16Fe0.04O4 +0.026FeMnO3
m4 800 13.230 177.265 0.308 Li1.03Mn1.72Ni0.16Fe0.04O4 *m5 875 13.104 175.577 0.126 Li0.8Mn1.72Ni0.16Fe0.04O4
*Pure phase identified from XRD patterns.
as P500 and P800 powders. The size distribution measurements are presented in Figures 4a and 4b. The size
distribution of the P500 powder (Figure 4a) exhibits a single population with an average size of 262 nm. The
size distribution from the P800 powder (Figure 4b) exhibits two populations near 220 nm and 2669 nm with an
average size of 274 nm. Comparison of the two distributions shows that raising the calcination temperature to
800 ◦C leads to formation of larger particles as the degree of crystallinity increases with the rising temperature.
The lower size population of the P800 powder is thought to arise from the fragments of the larger particles
due to the milling process. The powders’ morphologies, examined under SEM, are shown in the SEM images
of Figures 5a and 5b for the P500 and P800 powders, respectively. It is clear from the two SEM images that
P800 exhibits larger particles than P500. The P800 powder is less populated with small particles that are
produced from the milling effect than the P500 powder. Particle geometries of the P800 powder are also closer
to the octahedral shape of the spinel structure where milling conditions are less effective on highly crystalline
particles. Particles sizes were measured from the SEM images as shown by the arrows on each image. The sizes
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AL-TABBAKH et al./Turk J Phys
obtained from the SEM measurements were lower than those obtained from the particle size analyser. This is
justified by the fact that the size measurements using the particle size analyser are diffraction-limited, whereas
such limitations using electrons are insignificant in the SEM measurements [14]. Furthermore, aggregation of
particles may result in higher size measurements using the particle size analyser.
Size (nm) Size (nm)
P500
Inte
nsi
ty (
%)
(a)
0
5
10
15
20
25
30
164
190
220
255
295
342
396
459
1106
1484
3580
4145
1718
1990
2305
2669
3091
4801
1281
P800
Inte
nsi
ty (
%)
(b )
0
5
10
15
20
25
164
190
220
255
29
5
342
396
459
1106
1484
3580
4145
1718
1990
2305
2669
3091
4801
1281
Figure 4. The particle size distributions obtained for the calcined powders: (a) calcination at 500 ◦C, (b) calcination
at 800 ◦C. All powders were subjected to high-energy ball milling prior to measurements.
(b) (a)
Figure 5. Morphologies of powders as shown under scanning electron microscope: (a) powder calcined at 500 ◦C, (b)
powder calcined at 800 ◦C.
The XRD patterns of the two powders are presented in Figure 6. Peaks of high intensities were obtained
for both powders using the solid-state detector with a reasonable step-time. The high intensity peaks indicate
that highly crystalline powders were produced successfully by subjecting powders to the calcination conditions
(i.e. 500 ◦C or higher for 24 h). Intensities were found to increase with the rising of calcination temperature,
whereas the full-width at half-maximum (FWHM) of the main reflections decreased. Comparison of peaks
heights and FWHM are presented in Table 2. This shows that increasing the calcination temperature improved
the crystalline quality of the synthesised powders [16].
For the P800 powder, these peaks could be indexed to a single-phase of cubic spinel structure with space
group Fd3m (space group number 227) of lithium manganese oxide (ICSD 088138) [17]. No significant peaks
could be assigned to the impurity phase in the calcined powder. Unlike P800, the XRD patterns of P500
showed impurity peaks. The XRD patterns could be indexed with reference to two phases of a cubic spinel
structure (space group Fd3m) of lithium manganese oxide (ICSD 090132) and a cubic structure (space group
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AL-TABBAKH et al./Turk J Phys
2 �eta (degree)
Inte
nsi
ty (
cou
nts
)
0
100
400
900
1600
2500
3600
20 30 40 50 60 70 80 90
111
11
2
222
113
22
2
004
133
044
11
5
60
333
044
135
335
22
6
444
P500 P800
Figure 6. Comparison of XRD patterns of the P500 and P800 powders. Peaks indexed in black are common between
the two patterns.
Table 2. Peak parameters of main reflections of XRD patterns obtained from P500 and P800 powders.
Intensity (counts) FWHM (◦2θ)Peak (111) (113) (004) (111) (113) (004)P500 1158 424 467 0.44 0.56 0.65P800 3734 1095 1040 0.17 0.28 0.33
Ia3) of FeMnO3 (ICSD 1011266) [18,19]. XRD data analysis and Rietveld refinement suggested that P500 was
composed of 97% Li1.03Mn1.72Ni0.16Fe0.04O4 and 3% FeMnO3 as estimated from the intensity ratio of these
two phases (see Figures 7a–7c). This was found to agree with the TG analysis of the as-synthesised precursors
(i.e. 2.2% for the FeMnO3 phase), which was increased to about 2.8% after discarding residual water and
hydrocarbon impurities. The structural and stoichiometric parameters for the two powders obtained from the
refinement of the XRD patterns are presented in Tables 3 and 4.
2 �eta (degree)
(a)
(b)
(c)
20 30 40 50 60 70 80 90
111
11
1
112
222
22
2
112
113
11
3
004
00
4 00
4
004
114
113
222
004
23
3
134
044
226
22
613
6 22
614
5
224
044
115
224
115
044
135
Figure 7. Indexing of the main reflections of the XRD spectrum (spectrum (a)) obtained from the P500 powder to
Fd3m spinel phase (i.e. (b) ICSD: 090132, Fd-3m, LiMn2O4) and Ia3 cubic phase (i.e. (c) ICSD:1011266, Ia-3,
FeMnO3) .
The ionic distribution and the corresponding occupancies are shown in the same tables, where Mn, Ni, and
Fe ions occupy the octahedral 16d sites of the spinel structure with occupancies appropriate to the stoichiometric
composition of the spinel phase. Li ions occupy the tetrahedral 8a sites in addition to a fractional occupancy at
the octahedral 16d sites. The oxygen ions reside at the 32e sites of the unit cell forming the three-dimensional
frame of the spinel structure. This is in agreement with the reported results of the spinel structure, which has a
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AL-TABBAKH et al./Turk J Phys
Table 3. Structural and stoichiometric parameters of the P800 powder.
Formula of refined structure a (nm) Density (g/cm3) Rwp% GOFLi1.03Mn1.719Ni0.159 Fe0.04O3.98 0.820254(9) 4.245(4) 4.222 1.1443Space group (no.) and structure Fd3m (1) cubicAtom Wyckoff’s SOF X Y Z B(iso.)Li-a 8a 1.000 0.125 0.125 0.125 0.5439Li-b 16d 0.0149 0.500 0.500 0.500 0.5877Mn 16d 0.8596 0.500 0.500 0.500 0.4561Ni 16d 0.0798 0.500 0.500 0.500 0.4760Fe 16d 0.0200 0.500 0.500 0.500 0.5211O 32e 0.9958 0.25899 0.25899 0.25899 0.9083
Table 4. Structural and stoichiometric parameters of the P500 powder.
Formula Weight
fraction
Space-group no.
and structure a (nm)
Density
(g/cm3)
Rwp% GOF
Li1.0298Mn1.717Ni0.158Fe0.0396O3.943 97% la 3m (227)
cubic 0.817837 4.3035 2.2554 1.0382
Atom Wyckoff’s SOF X Y Z B (iso.)
Li-a 8a 1.000 0.125 0.125 0.125 0.51104
Li-b 16d 0.0149 0.500 0.500 0.500 0.52231
Mn 16d 0.8586 0.500 0.500 0.500 0.38484
Ni 16d 0.0793 0.500 0.500 0.500 0.43297
Fe 16d 0.0198 0.500 0.500 0.500 0.58013
O 32e 0.9858 0.261 0.261 0.261 0.77490
Formula Weight
fraction
Space-group no.
and structure a (nm)
Density
(g/cm3)
Rwp% GOF
FeMnO3 3 % la 3 (206) cubic 0.936530 5.1351 2.2554 1.0382
Atom Wyckoff’s SOF X Y Z B (iso.)
Fe-a 8a 0.5 0.00 0.00 0.00 0.5
Mn-a 8a 0.5 0.00 0.00 0.00 0.5
Fe-a 24d 0.5 0.28 0.00 0.25 0.5
Mn-b 24d 0.5 0.28 0.00 0.25 0.5
O 48e 1 0.28 0.365 0.13 0.5
unique character such that the manganese ions can easily be replaced with transition metal ions at the octahedral
sites [2,20]. The intensity of the 022 line is negligibly small in the XRD pattern, signifying that no transition
metal ions occupy the tetrahedral sites in the spinel phase. As additional proof, the ratio of the 004/113 lines
showed no evident decrease, which is in agreement with the previously reported studies for spinel structures
[3,20,21]. The site occupancy factor (SOF) of Li at the tetrahedral 8a site is full. At the octahedral sites,
fractional occupancy of Li ions, substituting some transition metals ions, is still possible. Similar findings have
been reported for nonstoichiometric spinel structures [3]. The unit-cell constant, obtained from the refinement
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AL-TABBAKH et al./Turk J Phys
results, shows an increase in value with the increase of calcination temperature. This suggests that formation of
the spinel phase is more complete at 800 ◦C than at 500 ◦C [22]. Furthermore, inspection of the site occupancy
of the constituting ions justifies the increase of the lattice parameter with temperature, as ion SOFs of P800 are
generally higher than the SOFs of P500. Increase of lattice parameter with temperature for the spinel structure
was reported by Ghanbarnezhad et al. and Yi et al. [23,24]. The potential profiles of the fabricated batteries
are shown in Figure 8. The open-circuit voltage decreases from 3.57 V for the P500 powder to 3.13 V for the
P800 powder with respect to Li anode potential, as shown in the inset of Figure 8. This could be attributed to
the effect of calcination temperature on the completeness of the spinel phase formation.
Po
ten
tial
(V
)
X (C/C Τ )
X (C/C T )
Po
ten
tial
(V
)
0.0 0 .5 1.0 1.5 2.0
1.0
1.5
2.0
2.5
3.0
3.5
4.0
4.5
0.0
X (C/C
0.01
C/C )C/C
0.02 0.03 0.03.0
Po
t
3.2
3.4 enti
al (
V3.6
(V
) 3.8
4.0
P500
P800
Figure 8. Potential profiles of the lithium ion batteries fabricated from P500 and P800 powders showing typical potential
plateaus of spinel structure. Inset: Charge curves at magnified scales showing open circuit potentials of the P500 and
P800 batteries.
Both batteries were initially charged to 4.2 V at 1 mA current. Cells were then discharged to 1.0 V at 0.5
mA to determine the equilibrium voltage profile. Above 3 V both cells exhibited a single potential plateau at
around 4.0 V, with capacity increases with the increase of calcination temperature. At 3 V potential, the P500
cell delivered 24% of its theoretical capacity, whereas the P800 cell delivered 43% of its theoretical capacity.
For further elaboration of this observation, the Li content dependence on the applied voltage was determined
from the SOF of Li in the tetrahedral sites of the unit-cell of the cathode active compounds (see Tables 3 and
4). Both powders exhibited similar SOFs of Li, signifying no deficiency in Li occupancy in the tetrahedral sites.
However, charging batteries to 4.2 V resulted in extraction of 24% of the Li from the tetrahedral sites of the
P500 compound and 43% of the Li from the P800 compound. It seems easier to extract more Li ions from
the active material when calcination temperature is higher (i.e. 800 ◦C). When the batteries discharged the
remaining amount of Li in each case has to be taken into consideration while estimating the amount of Li being
intercalated back into the active material (i.e. 76% of Li for P500 and 57% for P800). At 2.8 V (onset of the
second plateau), the Li content per formula in P500 and P800 was found equal to 1.14 and 0.99, respectively.
This implies that while 99% of the extracted lithium was intercalated back into the structure of the P800
cathode compound, overlithiation took place for the P500 compound. Due to further discharge to 2.5 V, the
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AL-TABBAKH et al./Turk J Phys
lithium content was found equal to 1.9 and 1.64 for the P500 and P800 active compounds, respectively, which is
determined from the discharge curves of Figure 8 after taking into account the remaining amount of Li content
at the cut-off voltage during battery charge. The potentials of the second plateau were found independent of
calcination temperature.
It is obvious from the potential profiles of the compounds that while the high-voltage plateau shows a
prominent linear variation of potential in accordance with the intercalation of lithium (i.e. a capacitor-like
behaviour), the second plateau exhibits a remarkably flat characteristic. For the P800 compound, the voltage
drop across the potential plateau is equal to 18.8 mV/h at 40 mA/g current density and increases linearly with
the discharge current. The P500 compound exhibits a slightly higher voltage drop across its potential plateau.
The capacitor-like behaviour has been attributed to low electronic conductivity and low lithium ion diffusion
coefficient [1] and, therefore, the usefulness of the battery material for industrial and the technological purposes
depends largely on moderating this problem. Based on that, it is proposed to operate the present batteries in
the voltage range of the low potential plateau where such capacitor-like behaviour is insignificant.
The charge and discharge curves of the first, second, and twentieth cycles in the voltage range of 3.75 V
to 2.5 V are shown in Figures 9a and 9b for the batteries fabricated from P500 and P800 powders, respectively.
Batteries were charged and discharged at 1 mA current. The deintercalation potentials (i.e. voltage of charge
plateau) of the P500 and P800 compounds were equal to 2.94 V and 2.95 V with respect to Li anode potential.
The voltages of the discharge plateaus were equal to 2.78 V and 2.82 V for the P500 and P800 batteries,
respectively. The charge and discharge potential plateaus increased with calcination temperature. This might
be attributed to the effect of calcination temperature on the formation of the spinel phase. The discharge-
specific capacities of the batteries over 20 cycles are shown in Figure 10. It exhibits a significant fade of 15%
of the initial capacity for the P500 batteries and an improvement of capacity of about 10% over cycling for
the P800 batteries. The attainable capacities after 20 cycles are 83.4 mAh/g and 96.5 mAh/g for the P500
and P800 compounds, respectively. The effect of calcination temperature on the coulombic efficiency is obvious
from Figure 10. The fact that the P500 compound exhibits a significant capacity fade over cycling may also
Specific capacity (mAh/g)
Vo
ltag
e (V
)
( b) P 8 00
0 20 40 60 80 100
1.0
4.5
4.0
3.5
3.0
2.5
2.0
1.5
20th 2nd
1st
Specific capacity (mAh/g)
Vo
ltag
e (V
)
(a) P500
0 20 40 60 80 100 1.0
4.5
4.0
3.5
3.0
2.5
2.0
1.5
20th
2nd 2nd1st
Figure 9. Galvanostatic charge and discharge curves of the first, second, and twentieth cycles of the batteries: (a)
battery fabricated from P500 powder, (b) battery fabricated from P800 powder.
10
AL-TABBAKH et al./Turk J Phys
be understood based on analysis of the potential profile of Figure 8, where this material was found to undergo
overlithiation at relatively lower voltages than the P800 compound. Overlithiation of the spinel structure is
known to cause irreversible structural changes that finally lead to loss of capacity [25]. The dependence of
specific capacity on discharge current for the P800 compound was investigated as shown in Figure 11, which
exhibits a linear behaviour. The discharge capacity decreased from 94.6 mAh/g at 1 mA (i.e. current density
of 40 mA/g) to 74 mAh/g at 2.5 mA (100 mA/g).
Cycle (n)
Dis
char
ge-s
pec
ific
cap
acit
y (m
Ah
/g)
P500
0 2 4 6 8 10
50
110
100
90
80
70
60
12 14 16 18 20
P800
Discharge current (mA)
Spec
ific
cap
acit
y (m
Ah
/g)
0.5 1.0 1.5
70
100
95
90
85
80
75
2.0 2.5
Figure 10. Discharge-specific capacities of the P500 and
P800 batteries over 20 cycles.
Figure 11. The dependence of specific capacity on dis-
charge current for the P800 battery.
4. Conclusions
From the present analysis, a cathode compound of the form Li1.03Mn1.72Ni0.16Fe0.04O4 was successfully
synthesised for lithium rechargeable battery application. The lithium ion batteries, fabricated from the present
powders, exhibited outstanding features of flat potential plateau and minimal voltage drop, emphasising a
minimal capacitor-like behaviour. The effect of calcination temperature on battery performance was such that
batteries prepared from powders calcined at 800 ◦C exhibited higher coulombic efficiency and better cycling
performance than powders calcined at 500 ◦C. These features were attributed to elimination of the impurity
phase and development of a single phase structure. The present material is thought to be a potential candidate
for energy storage applications as a high-rate rechargeable battery.
Acknowledgment
One of the authors, AAA Al-Tabbakh, thanks the Universiti Teknologi MARA for support through a postdoc-
toral fellowship.
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